PL208233B1 - Cold rolled metal plate of 780MPa or more tensile strength, ideal selectional compliance and a decrease in weld hardness - Google Patents

Abstract

The present invention provides a high-strength cold-rolled steel sheet and a high-strength surface treated steel sheet 780 MPa or more in tensile strength, said steel sheets having excellent local formability and suppressed weld hardness increase and being characterized by: said steel sheets containing, in weight, C: 0.05 to 0.09%, Si: 0.4 to 1.3%, Mn: 2.5 to 3.2%, P: 0.001 to 0.05%, N: 0.0005 to 0.006%, Al: 0.005 to 0.1%, Ti: 0.001 to 0.045%, and S in the range stipulated by the following expression (A), with the balance consisting of Fe and unavoidable impurities; the microstructures of said steel sheets being composed of bainite of 7% or more in terms of area percentage and the balance consisting of one or more of ferrite, martensite, tempered martensite and retained austenite; and said components in said steel sheets satisfying the following expressions (C) and (D) when Mneq. is defined by the following expression (B); S≰0.08×(Ti(%)−3.43×N(%)+0.004 . . . (A), where, when a value of the member Ti(%)−3.43×N(%) of said expression (A) is negative, the value is regarded as zero. Mneq.=Mn(%)−0.29×Si(%)+6.24×C(%) . . . (B), 950≰(Mneq./(C(%)−(Si(%)/75)))×bainite area percentage (%) . . . (C), C(%)+(Si(%)/20)+(Mn(%)/18)50.30 . . . (D).

Description

Description of the invention

The present invention relates to cold-rolled steel sheet, in particular surface treated.

Until now, steel sheets with a tensile strength of 590MPa or less have been used for parts that mainly make up the body of cars or motorcycles. In recent years, work has been done to increase the strength of the material to a higher level, and additionally to use steel sheets with increased strength, in order to reduce the weight of the car body and improve fuel efficiency and improve safety in the event of a collision.

High-strength steel sheets manufactured to meet the above purposes are mainly used for car body frame members and reinforcement members, seat frame parts, and steel sheet with a tensile strength of at least 780MPa and having excellent formability is especially desirable.

Such parts are processed such as forming by stamping and rolling. However, due to the demands of body designers and other industrial designers, it is sometimes difficult to drastically change the shape of such parts to which a conventional steel sheet is used with a tensile strength of 590MPa or less and therefore, to allow complex shapes to be shaped high-strength sheet steel having excellent workability is required.

In the meantime, the processing methods have changed from the usual extrusion with a half-shape fixture to simple stamping or bending processing suitable for use with high-strength steel plate. In particular, when the bending edge curves into an arc of a circle or the like, sometimes the ends of the steel sheet elongate, i.e. a stretched flange is formed. In addition, a deburring treatment is often used for some parts, where the collar is formed by widening the treatment hole (bottom hole). In some high magnification cases, the orifice is widened to a diameter that is at least 1.6 or more of the exit diameter. At the same time, the phenomenon of elastic recovery after machining the part, such as elastic shrinkage, occurs as the strength of the steel plate increases and harms the maintenance of the accuracy of the part. For these reasons, the idea of reducing the inside bend radius, for example to about 0.5 mm, is often used in forming processes, especially in bending machining.

However, in such processing, although the steel sheet is required to have local formability such as stretched flange formability, hole expandability, flexibility and the like, typical high strength steel sheet is insufficient to provide such formability and therefore the problem of typical high strength steel poses problems such as cracking. and the product cannot be processed stably.

At the same time, such stamped parts are often joined to other parts by spot welding or other welding. However, in the case of high-strength steel sheet with a tensile strength of at least 780MPa, a metallurgical method of increasing the C content of the steel is often used as a means of effectively ensuring strength, and the problem with using such a method is that the welded metal is extremely hardened by heating and cooling during welding, and therefore the properties of the weld and the product are deteriorated.

A hitherto known high strength steel sheet having improved deformability of a stretched flange is that proposed in the publication of the unexamined Japanese Patent Application No. H9-67645. However, this technology only improves the stretched flange formability after punching and does not necessarily improve the properties of the weld.

In addition, Japanese Patent Nos. H2-1894 and H5-72460, tested, propose methods for improving the weldability of high-strength steel sheets. This prior technology improves the cold workability and weldability of the high-strength steel plate. However, with regard to improving the cold workability of this technology, the improvement of local formability such as stretched flange shaping, hole expandability, flexibility and the like is not sufficiently proven. In contrast, the latter technology presents an improvement in the formability of the stretched flange and, in addition, the weldability. However, the strength of the steel sheet covered by the solution is around 550MPa, and this technology does not apply to high strength steel sheet with a tensile strength of at least 780MPa.

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In addition, research has led to the following statements. In the case of a high-strength starting steel sheet with a tensile strength of at least 780MPa, the main tensile mechanism is primarily responsible for the hard martensite and bainite in the second phase, and the C content of the steel acts as the main factor of the reinforcing mechanism. However, as the C content increases, the local formability probably deteriorates, and at the same time the hardness of the weld visibly increases. However, due to the above-mentioned problems of a high strength steel sheet of a starting steel with a tensile strength of at least 780MPa, no suggestions have been found for improving local formability and reducing weld hardening.

According to the invention, a cold-rolled steel sheet, in particular surface treated, with a tensile strength of at least 780MPa, is characterized in that it contains (in% by weight):

C: 0.05 to 0.09%,

Si: 0.4 to 1.3%,

Mn: 2.5 to 3.2%,

P: 0.001 to 0.05%,

N: 0.0005 to 0.006%,

Al: 0.005 to 0.1%,

Ti: 0.001 to 0.045% i

S: within the range specified by the following expression (A),

S <0.08 x (Ti (%) - 3.43 x N (%)) + 0.004 (A), in which, when the value of the term Ti (%) - 3.43 x N (%) of the expression (A) is negative , the value is considered to be zero, the remainder being Fe and the unavoidable impurities, whereby the indicated components satisfy the following expressions (C) and (D):

950 <(Mneq./(C(%) (Si (%) / 75))) x percentage of bainite area (%) (C),

C (%) + (Si (%) / 20) + (Mn (%) / 18) <0.30 (D), where Mneq. is determined by the following expression (B):

Mneq. = Mn (%) - 0.29 x Si (%) + 6.24 x C (%) (B), and the microstructure of the steel sheet has a bainite area percentage of at least 7%, the remainder being at least one from ferrite, martensite, tempered martensite and residual austenite.

Preferably, the steel sheet comprises as additional components at least one of (in% by weight):

Nb: 0.001 to 0.04%

B: 0.0002 to 0.0015%, i

Mo: 0.05 to 0.50%.

Preferably, the steel sheet comprises 0.0003 to 0.01 wt.% Ca as an additional component.

Preferably, the steel sheet comprises 0.0002 to 0.01 wt.% Mg as an additional component.

Preferably, the steel sheet comprises 0.0002 to 0.01% by weight of rare earths as additional components.

Preferably, the steel sheet contains as additional components from 0.2 to 2.0% by weight Cu and from 0.05 to 2.0% by weight Ni.

The steel sheet is preferably coated with zinc or its alloy in the surface treatment.

Steel sheet according to the invention, cold rolled and high strength, especially surface treated, from a starting steel having a tensile strength of at least 780MPa, which steel sheet has excellent local formability such as stretched flange shaping, hole expandability, flexibility and the like, it is distinguished by a reduced increase in the hardness of the weld and good welding properties.

The subject matter of the invention is illustrated with reference to the drawing, in which Fig. 1 is a graph showing the effect of the value of the term to the right of the inequality sign in the expression (A), which defines the upper limit of the S content, and the effect of the S content on the local formability index, Fig. 2 is a graph showing the relationship between the value of the right-hand end of the inequality sign in expression (C) and the overall widening ratio as an indicator of local deformability, and fig. 3 is a graph showing the effect of the left-hand end of the inequality sign in expression (D ) to increase the hardness of the weld.

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The steel composition and metallographic structures of the steel sheets were investigated in terms of means of reducing the weld hardness while ensuring the local deformability of the steel sheet, such as stretched flange forming, hole expandability, flexibility and the like. First of all, as a result of the local deformability test of the steel sheet, it was found that in the case of high-strength steel sheet of the starting steel with a tensile strength of at least 780MPa, the local deformability is mainly determined by the form of the metallographic structure of the steel sheet and the ease of formation of inclusions in it, such as like extractions and the like. Moreover, it has been found that the local formability can be improved by the content of: C, Si, Mn, P, S, N, Al, Ti, of which of these components S, Ti and N act as factors determining the formation of sulfide-type inclusions satisfying the given expression determining the relationship, and further by regulating not only the content range of a particular component, such as C, but also the relationship between a structure favorable to local formability and a plurality of components, including C, acting as a cure index.

In the production of steel sheet having a tensile strength of at least 780MPa, measures employing a hardened structure of martensite, bainite and the like have generally been used. For example, it is widely known that in the case of a steel sheet of the two-phase complex structure type (two-phase steel sheet), excellent plasticity, near the interface between the soft ferrite phase and the hard martensite phase, many dislocations capable of dislocation, formed by cooling, are introduced near the interface between the soft ferrite phase and the hard martensite phase. and therefore a high elongation is obtained. However, the problem of such a steel sheet is that the microscopic structure is inhomogeneous due to the coexistence of the soft and hard phases, and as a result, the hardness difference between the phases is large, and the interface between the phases cannot withstand the local deformation and cracks are formed. Therefore, in order to solve the problem, it is effective to homogenize the structure in the case of a single phase martensite structure, bainite structure or tempered martensite structure. In particular, the bainite structure is excellent at creating a balance between strength and plasticity, and shows good workability. In light of the above facts, it was found that the ease of obtaining the required bainite structure is strongly influenced by C, Si and Mn, and the local formability is improved when these elements and the actual percentage of the obtained bainite structure satisfy the given expression determining the relationship.

In addition, as a result of research on preventing the increase in weld hardness, it has been found that the increase in hardness is due to the martensitic transformation that occurs with rapid cooling after rapid local heating during welding, and the increase in the hardness of the weld is effectively reduced when C and Si, affecting hardening, Mn is satisfied by the given dependency expression.

The present invention will be explained in more detail below.

First, the reasons for controlling the components of the steel are explained further.

C is an element important for increasing the strength and hardness of the steel, and is essential for obtaining a complex structure consisting of ferrite, martensite, bainite, and the like. In particular, an amount of 0.05% or more of C is necessary to provide a tensile strength of at least 780MPa and an effective amount of bainite structure favorable for local formability. On the other hand, if the C content increases, not only is it difficult to obtain a bainite structure, but iron carbides such as cementite tend to grow and consequently the local formability deteriorates and the hardness noticeably increases after welding, causing the weldability to deteriorate. . For these reasons, the upper limit of the C content is set at 0.09%.

Si is an element beneficial for increasing the strength without compromising the workability of the steel. However, when the Si content is less than 0.4%, a pearlite structure tends to deteriorate local formability, and the difference in hardness between the shaped structures increases due to the reduction in the strength of the ferrite by the dissolved element, which deteriorates the local formability. For these reasons, the lower limit of the Si content is set at 0.4%. On the other hand, when the Si content exceeds 1.3%, the cold rollability deteriorates due to the increase in the solubility of the ferrite in the hardening compounds, and the phosphate treatability deteriorates due to oxides formed on the surface of the steel sheet. The sealability also deteriorates. For these reasons, the upper limit of the Si content is set at 1.3%.

Mn is an element that effectively increases the strength and hardening of steel, and provides a bainite structure favorable for local formability. When the content of Mn is less than

PL 208 233 B1

2.5% a satisfactory structure is not obtained. Therefore, the lower limit of the Mn content is set at 2.5%. On the other hand, when the Mn content exceeds 3.2%, the workability of the base steel deteriorates and the weldability also deteriorates. For this reason, the upper limit of the Mn content is set at 3.2%.

A P content of less than 0.001% increases the cost of dephosphorization and therefore the lower limit of the P content is set at 0.001%. On the other hand, when the P content exceeds 0.05%, there is a lot of segregation during the solidification during casting, and therefore internal crack formation and workability deterioration. In addition, there is also brittleness of the weld. For these reasons, the upper limit of the P content is set at 0.05%.

S is an element that is extremely detrimental to local formability because it remains as sulfide inclusions such as MnS. In particular, the effect of S increases as the strength of the starting steel increases. Therefore, when the tensile strength is at least 780MPa, S should be limited to 0.004% or less. However, when Ti is added, the effect of S is mitigated to some extent as Ti separates as titanium sulfide. Therefore, in the present invention, the upper limit of the S content can be adjusted by the following expression (A) considering the Ti and N content:

S <0.08 x (Ti (%) - 3.43 x N (%)) + 0.004 (A), where, when the value of the term Ti (%) - 3.43 x N (%) of the expression (A) is negative, the value is considered zero.

Al is an element necessary for steel deoxidation. When the Al content is less than 0.005%, deoxidation is not sufficient, bubbles remain in the steel and defects such as pores are formed. Therefore, the lower limit of the Al content is set at 0.005%. On the other hand, when the Al content exceeds 0.1%, the occurrence of inclusions such as aluminum oxide increases and workability deteriorates. Therefore, the upper limit of the Al content is set at 0.1%.

Considering N, an N content of less than 0.0005% increases the cost of cleaning the steel. Therefore, the lower limit of the N content is set at 0.005%. On the other hand, when the N content exceeds 0.006%, the workability of the base steel deteriorates, the possibility of thick TiN formation when combining N and Ti is high, and local formability deteriorates. Moreover, hardly any Ti remains necessary for the formation of titanium sulphides, which is disadvantageous in reducing the upper limit of the S content proposed in the present invention. Therefore, the upper limit of the N content is set at 0.006%.

Ti is an element that effectively creates titanium sulphides, which have a relatively weak effect on local formability and reduce harmful MnS sulphides. Moreover, Ti reduces the grain growth of the weld metal structure and prevents its brittleness. Since a Ti content of less than 0.001% is insufficient to produce this effect, the lower limit of the Ti content is set at 0.001%. In contrast, when Ti is added excessively, not only does the size of the rectangular precipitates of TiN increase, which worsens the local formability, but also a stable carbide is formed, which lowers the C content in austenite during the production of the starting steel and does not allow to obtain the structure of the required hardness, and thus, it does not provide the desired tensile strength. For these reasons, the upper limit of the Ti content is set at 0.045%.

Nb is an element effectively forming fine carbides that reduce the softness of the heat affected zone of the weld and may be added. However, when the Nb content is less than 0.001%, the softening effect of the heat affected zone of the weld is not sufficient. Therefore, the lower limit of the Nb content is set at 0.001%. On the other hand, when excessive amounts of Nb are added, the workability of the starting steel is degraded by increasing carbides. Therefore, the upper limit of the Nb content is set at 0.04%.

B is an element effective for improving the hardening of the steel and reducing the diffusion of C in the heat-affected zone of the weld, and thus softens it due to the action of C, and can therefore be added. An addition of 0,0002% or more of B is necessary to produce this effect. On the other hand, when excess B is added, not only does the workability of the starting steel deteriorate due to the growth of carbides, but also makes the steel brittle and deteriorates when hot. For these reasons, the upper limit of the B content is set at 0.0015%.

Mo is an element that facilitates the formation of a bainite structure. Moreover, Mo has the effect of reducing the softness of the heat affected zone of the weld, and this effect is expected to increase further in the presence of Nb or the like. Therefore, Mo is a preferred element for improving the quality of the weld and can be added. However, the addition of Mo in an amount of less than 0.05%

PL 208 233 B1 is insufficient to produce these effects and therefore its lower limit is set at 0.05%. In contrast, even when Mo is added excessively, the effects saturate, and this is not economically advantageous. For these reasons, the upper limit of the Mo content is set at 0.50%.

Ca improves the local formability of the starting steel by adjusting the shape (spheroidization) of the sulphide inclusions and can be added. However, the addition of Ca less than 0.0003% is insufficient to produce this effect. Therefore, the lower limit of the Ca content is set at 0.0003%. On the other hand, even when Ca is added in excess, not only is the effect saturated, but also the opposite effect (deterioration of local formability) increases due to the increase in precipitates. Therefore, the upper limit of the Ca content is set at 0.01%. It is desirable that the Ca content be 0.0007% or more for a better effect.

Mg on addition forms oxides by combining with oxygen, and it is judged that the MgO so formed or the complex oxides of Al2O3, SiO2, MnO, Ti2O3 and the like containing MgO separate out as very fine phases. Although this is not confirmed, it is sufficiently estimated that the size of each precipitation is small and therefore the precipitates are statistically distributed uniformly. In addition, it is assessed, although not obvious, that such finely dispersive oxides, evenly distributed throughout the steel, create fine voids in the punch or shear plane, from which cracks begin during stamping or cutting, reduce stress concentrations during flanging and flange shaping processing. and hence prevent the growth of fine pores that increase cracks. Therefore, Mg may be added to improve the hole expandability and the deformability of the stretched flange. However, the addition of Mg less than 0.0002% is insufficient to achieve these effects and therefore the lower limit is set at 0.0002%. On the other hand, when the Mg addition exceeds 0.01%, not only is there no improvement in proportion to the amount added, but also the cleanliness of the steel and the hole expandability and the deformability of the elongated flange are deteriorated. For these reasons, the upper limit of the Mg content is set at 0.01%.

Rare earth metals (REM) are elements that have the same effect as Mg. While not sufficiently validated, it is judged that REM are elements that improve hole expansion and elongated flange formability by reducing fine oxide formation cracks and therefore REM may be added. However, when the REM content is less than 0.0002%, the effect is insufficient and therefore the lower limit is set at 0.0002%. On the other hand, when the REM addition exceeds 0.01%, not only is there no improvement in proportion to the amount added, but also the cleanliness of the steel and the hole expansion and the formability of the elongated flange are deteriorated. For these reasons, the upper limit of the REM content is set at 0.01%.

Cu is an element that efficiently improves the corrosion resistance and fatigue resistance of a base steel, and may be added as required. However, when the Cu content is less than 0.2%, the corrosion and fatigue resistance improvement effect is insufficient, and therefore the lower limit is set at 0.2%. On the other hand, excessive Cu addition saturates the effect and increases costs, and therefore its upper limit is set at 2.0%.

In steel with the addition of Cu, sometimes during hot rolling, surface defects, called copper scales, may form, due to insufficient heat. The addition of Ni prevents the formation of a copper flake, and when Cu is added, an amount of Ni is provided that is set at 0.05% or more. On the other hand, an excess of Ni addition saturates the effect and increases costs. Therefore, the upper limit of the Ni content is set at 2.0%. Here, the effect of adding Ni indicates that it is desirable that the amount of Ni added be such that the Ni / Cu weight ratio is in the range of 0.25 to 0.60.

For high-strength, cold-rolled steel sheets, hole expansion tests were carried out, the results of which were considered a typical indicator of local formability, and the relationship between the expression (A), which defines the upper limit of the S content, and the S content, was examined. The results are shown in Fig. 1. Obtained excellent local formability when the content of S is within the range defined by the expression (A). In Fig. 1, the mark O represents an opening expansion ratio greater than 60% and X represents an opening expansion ratio less than 60%. It can be seen from the graph that when the amounts of S, Ti and N additive are within the ranges defined by the present invention, the hole expansion ratio is 60% or greater, and the local formability is excellent.

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The above data show that at the upper limit of the S content, its influence to some extent is smaller due to the formation of titanium sulphides, which reduce the effect of MnS deteriorating local formability. This is a proposal different from the previous proposed methods where local formability is improved only by reducing the amount of S, and is also suitable for reducing the increase in desulfurization costs.

Moreover, in the present invention, the percentage area of the bainite structure and the amount of C, Si and Mn must satisfy the following expression (C):

Mneq. = Mn (%) - 0.29xSi (%) + 6.24xC (%) (B),

950 <(Mneq./(C(%) - (Si (%) / 75))) x area percentage of bainite (%) (C).

Through the above-mentioned experiments, the relationship between the value of the term on the right-hand side of the above expression (C) and the expansion ratio of the hole serving as an indicator of local formability was investigated. The results are shown in Fig. 2. In Fig. 2, the mark O represents an aperture expansion ratio greater than 60% and X represents an aperture expansion ratio less than 60%. It can be concluded from the diagram that when the state of the microstructure formed and the amounts of additive C, Si and Mn satisfy the expression relationship, the hole expansion ratio is 60% or greater, and the local formability is excellent.

The above data proves that when the value is related not only to the amount of bainite structure favorable for local formability, but also to hardening elements such as C, Si and Mn, the main influence on the formation of this structure is less than the value of the left term, and does not obtain sufficient local formability.

At the same time, according to the present invention, the amounts of C, Si and Mn must satisfy the following relationship (D):

C (%) + (Si (%) / 20) + (Mn (%) / 18) <0.30 (D).

In the above-mentioned tests, the relationship between the value obtained according to the above relation (D) and the maximum hardness of the spot welding weld and the shape of the fracture in the weld tensile test was investigated. The results are shown in Fig. 3. The horizontal axis shows the value calculated from the left-hand segment of expression (D), and the vertical axis shows the ratio of the maximum spot weld hardness to the base steel hardness (ratio of weld hardness to base steel - K), and each hardness was measured by the Vickers method (under a load of 100 gf) in a part of one fourth of the thickness of the sheet on the cross-sectional area. In Figure 3, the mark O represents the cases where the weld hardness to base steel ratio K is less than 1.47, and the X mark shows the cases where the weld hardness to base steel ratio K is greater than 1.47. It is understood from the figure that when the amounts of additive C, Si and Mn are in the control range according to the present invention, the increase in the hardness of the weld is suppressed and the ratio K is not greater than 1.47. At the same time, a breakthrough occurs in the weld nucleus when this K ratio exceeds 1.47, and when the K ratio is not greater than 1.47, the breakthrough occurs outside the weld nucleus, so the weldability is good.

The above-mentioned relation (D) determines the composition range for which the hardness of the martensite formed by cooling is reduced during heating and rapid cooling of the weld.

Moreover, auxiliary components, such as Cr and V and the like, inevitably included in the steel sheet, are not at all detrimental to the properties of the steel according to the invention. However, excessive addition of ingredients may increase the recrystallization temperature, deteriorate the rollability and also the workability of the starting steel. For these reasons, with regard to these auxiliary components, it is desirable to regulate the Cr content to 0.1% or less and V to 0.01% or less.

The method for producing a high-strength cold-rolled steel sheet, in particular a high-strength surface treated steel sheet according to the invention, can be properly selected according to the application and the required properties.

In the present invention, the above-indicated components form the basis of the steel according to the invention. When the bainite area percentage is less than 7% of the microstructure of the base steel, local formability hardly improves. Therefore, the lower limit of the bainite area percentage is set at 7%. A preferred bainite area percentage is 25% or greater. The upper limit of the bainite area percentage is not particularly fixed. However, when it exceeds 90%, the ductility of the base steel deteriorates due to an increase in the hard phase and the use of press parts is severely limited. Therefore, the preferred upper limit of the bainite area percentage is set to 90%. At the same time, the influence of a different microstructure on the workability of the base steel must be taken into account and, in order to secure the balance

Between workability and ductility, the preferred ferrite area percentage is at least 4%.

Steel adapted to contain the above-mentioned components was processed by the following example method and steel sheets were produced. First, the steel was smelted and refined in a converter and cast into slabs by a continuous casting process. The resulting slabs were placed in a heated furnace at their high temperature, or, after cooling to room temperature, heated to a temperature in the range of 1150 ° C to 1250 ° C, and then subjected to a final rolling temperature in the range of 800 ° C to 950 ° C. C and coiled at a temperature of 700 ° C or less, and finally hot rolled steel sheets were produced. When the final temperature is less than 800 ° C, the crystalline grains are in a mixed grain state and the workability of the starting steel deteriorates. On the other hand, when the final temperature exceeds 950 ° C, the austenite grains are coarser and it is difficult to obtain the desired microstructure. A coiling temperature of at most 700 ° C is acceptable. However, at lower temperatures there is a tendency to inhibit pearlitic structure formation and the microstructure envisaged in the present invention can be obtained. Therefore, a preferred coiling temperature is 600 ° C or less.

The hot rolled steel sheets were then pickled, cold rolled and then annealed to produce cold rolled steel sheets. Although the reduction ratio in cold rolling is not particularly envisaged, an industrially preferred range is from 20 to 80%. The annealing temperature is essential to ensure the anticipated strength and workability of the high strength steel sheet, and its preferred range is from 700 ° C to below 900 ° C. When the annealing temperature is lower than 700 ° C, insufficient recrystallization occurs and it is difficult to obtain a stable workability of the starting steel. On the other hand, when the annealing temperature is 900 ° C or higher, the austenite grains are coarser and the required microstructure cannot be obtained. Moreover, a continuous annealing process is advantageous to obtain the microstructure envisioned by the present invention. In the case of high strength surface treated steel sheet, the galvanization of the cold-rolled steel sheet produced according to the above method is used under conditions in which the steel sheet is heated below 200 ° C.

For example, when galvanizing is used, the surfaces of the steel sheet are coated with 3 mg / m ² to 80 g / m ² . When the amount of coating is less than 3 mg / m ^{2, the} protective function of the coating against corrosion is insufficient and the purpose of electroplating is not fulfilled. On the other hand, when the coating amount exceeds 80 g / m ^{2, the} economic performance is deteriorated and there is a high tendency for defects such as blisters to form during welding. For these reasons, the preferred coating amount range is within the limits indicated above.

Moreover, even when an organic or inorganic coating is applied to the surface of a cold-rolled steel or plating layer, the effects of the present invention are not null and void. It should be noted that also in this case the temperature of the steel sheet should not exceed 200 ° C.

In this way, a high-strength cold-rolled steel plate is obtained, in particular a high-strength surface-treated steel plate, with a tensile strength of 780MPa or more, with excellent local formability and reduced weld hardness increase.

Examples

Steels containing the components shown in Table 1 were melted and refined in a converter and cast into slabs by a continuous casting process. Then the resulting slabs were heated to a temperature in the range of 1200 ° C to 1240 ° C, then hot rolled at a final temperature in the range of 880 ° C to 920 ° C (sheet thickness: 2.3 mm) and rolled at a temperature of 550 ° C or less. The produced hot rolled steel sheets were then cold rolled (sheet thickness: 1.2mm), heated respectively to a prescribed temperature ranging from 750 ° C to less than 880 ° C in a continuous annealing process, and then subjected to slow cooling to a predetermined temperature. temperatures ranging from 700 ° C to 550 ° C, and subsequent cooling thereafter.

The high-strength steel sheets produced in the above experiments were tested in the rolling direction and perpendicular to the rolling direction with JIS # 5 test specimens. The hole expansion ratio was then measured according to the hole expansion test method provided by the Japan Iron and Steel Federation Standard. Then, the percentage of bainite area was measured on the sections in the direction of rolling of steel sheets by carrying out the following steps: subjecting the sections to mirror grinding, subjecting them to

Corrosion treatment for highlighting residual γ etching (Nippon Steel Corporation, Haze: CAMP-ISIJ, vol. 6 (1993), p. 1698), observing the microstructure under 1000 magnification in an optical microscope, and applying image processing. The percentage of bainite area was defined as the mean value observed over the ten observation fields in terms of dispersion.

Moreover, for these high-strength steel sheets, spot welding was used for certain grades of high-strength steel sheets and the weld was assessed. Spot welding was carried out under spatter-free conditions using a 6 mm diameter dome electrode with a pressure of 400 kg and a nucleus diameter greater than four times the square root of the sheet thickness.

The weld was assessed by a shear strength test.

As for the increase in the hardness of the weld, the hardness was measured by the Vickers method (measurement load: 100 gf) in intervals of 0.1 mm in a part of one quarter of the sheet thickness on the cross-sectional area containing the weld, and the ratio of the maximum hardness of the weld to the hardness of the starting steel was measured, and the quality of the weld was then assessed. The results are shown in Table 2.

It should be concluded from the table that the steels according to the invention have excellent local formability and a inhibited increase in the hardness of the weld with respect to the reference steels.

The present invention provides a high strength cold rolled steel plate, especially a high strength surface treated steel plate with a tensile strength of 780MPa or more, with excellent local formability and reduced weld hardness increase.

PL 208 233 B1

Table 1

Comments Steel according to the invention | - Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Steel according to the invention Expression D en θ ' r ~~ CM about 0.29 0.25 0.27 no cm about in ΓΜ about o. o CM about about CM ^ about 0.26 0.26 C4 about 0.26 0.26 0.26 MD CM ^ about 0.25 Expression B about o. o cm cm 3.24 3.01 2.86 2.90 m 3.49 3.49 3.00 2.68 2.97 3.06 2.93 2.93 sD r — M en 3.04 Expression AND 0.0054 0.0056 0.0040 0.0040 0.0051 0.0040 0.0064 0.0040 0.0062 0.0065 0.0046 0.0042 0.0040 0.0040 0.0040 0.0040 0.0040 chemical substances of steel (wt%) Other ingredients chemical 1 1 1 1 1 1 1 1 vol 1 B: 0.0007 CM © © -D WITH about CM, © about 2 oo CM CM O · - about § ?: ® 2 Cd UJ u 2 CM CM © θ ' about eh 2 Cu: 0.046 Ni: 0.24 H. 0.025 ABOUT\ en ABOUT about 0.006 0.001 0.029 0.018 00 m CD about 0.003 0.041 about en about 0.015 ΟΊ ABOUT about 0.004 0.009 0.006 ABOUT θ ' about 0.005 WITH 0.002 0.006 0.005 en about about θ ' 0.004 un ABOUT about_ about 0.002 0.003 0.041 ABOUT\ en about 0.015 3 _λ cont no about about 0.009 0.006 0.001 0.005 Al 0.042 ir » m ABOUT about 0.042 0.037 0.028 sO en ABOUT about 0.057 0.029 0.094 j 0.038 0.034 0.044 And 0.037 0.045 0.047 0.041 0.035 0.0050 0.0052 0.0005 0.0024 0.0049 0.0037 0.0028 0.0027 0.0028 0.0059 0.0018 0.0027 0.0032 0.0033 CM © θ ' about 0.0032 0.0033 cu no'o 0.015 ' 0.014 0.010 0.009 0.007 r- about ABOUT about 0.008 ί 0.012 0.022 0.047 0.012 0.011 0.010 0.008 0.009 0.015 are 2 those CM 3.1 cm ογ cm ABOUT; cm Ch cm m m MD cm cm © en cm cm o. o cm 00 cm so cm <75 0.44 1.25 ογ Cd 0.47 1.16 0.51 about 0.43 about so about 0.56 ^ r 60Ί UD about 0.72 0.77 0.57 0.40 Composition AT 0.06 he ο _λ o 0.07 about about 0.05 0.06 0.06 about about about Ch about about 0.08 0.05 0.05 0.09 o. o © about 0.07 o. o θ ' about' 0.09 Steel code < m AT Q here ABOUT X 2 with ABOUT CU σ

1) Numbers in the shaded boxes are outside the ranges provided for by the invention.

PL 208 233 B1

PL 208 233 B1

Comments Steel according to the invention | Steel according to the invention | Sial according to the invention | 1 1 > £ ac .2 about 4> £ 3 00 Steel according to the invention Stable according to the invention | Steel according to the invention | and Steel according to the invention | 1 Steel according to the invention | t Steel according to the invention j | He stood according to the invention Steel according to the invention | Steel according to the invention ( Steel according to the invention I Steel according to the invention | He stood according to the invention 1 Steel according to the invention | Shape breakthrough welds point ! outside the testicle ( £ g Ϊ about. E. 2 and .5 2 Ί & Ϊ Ί and outside the j nucleus outside the kernel | S. 2 -and S. and 'outside the kernel | and • -! s and outside the nucleus j | outside the kernel | 1 outside the testicle] [a suitor | outside the nucleus | outside the nucleus S. 2 s Ϊ © about. 1 outside the kernel! λ 73 3 3 ll 3 ABOUT ABOUT about about ABOUT ABOUT about ABOUT about ABOUT ABOUT ABOUT ABOUT ABOUT ABOUT ABOUT ABOUT K ratio of the hardness of the weld - steel of the outputs, (K ^s max. Hard, welds / max. Hardness of the steel of the outputs,) Ok cm Ok CM ^ 3 m about CM © 3 Ok en 3 © cm CM Ck CM «Ck en © CM eo CM rc £ §% · § o and § a> δ CM r— en CS r © ABOUT 3 00 3 o. o »N m «Η Ok m en ABOUT 3 © CM 3 Ok CM 3 Axis about 3 3 Vk <3 © o. o en en 3 ABOUT ABOUT 3 Ok Ok en C- 3 <Λ 3 Hardness steel out- cove (HV 0.1) Axis o © CM Ok r * CM ABOUT en Ok 3 en en about 3 m © «Λ en thirty •ABOUT en Ok Ok CM Vk CM en CM Ok ΓΜ 3 en ί- ο «3 C- en Ok r * k en <n 3 m C- m Rating place- in deform- love λ> 60% ABOUT ABOUT ABOUT ABOUT ABOUT ABOUT ABOUT about ABOUT ABOUT ABOUT ABOUT about ABOUT ABOUT ABOUT ABOUT Hole expansion ratio λ (%) CM r- CM Ok ABOUT CM and ~ - Ok r- o. o r- about OO r- r- 00 r- 3 © Ok ABOUT Ok 00 Ok © thirty 3 •thirty o. o £ c * Tensile strength (MPa) CM Ok 3 tA Ok c- ABOUT 00 ag ABOUT * n Ok ί 1054 1 I 1077 I 3 CM 3 about CM 3 Ok 3 CM in | 1005 en Ok Ok | 1005 in k © c © k © o. o © CM Ok Expression C. r- © about CM about 00 Wk CM thirty 3 CM 00 c- 00 13053 Ί en 3 Ok <· *> en CM c- o. o 3 δ © c- en 14668 k © k © en 3 o. o Γ- CM • Π about m o. o • Zk CM © § Expression D CA <μ _λ θ ' r- CM ABOUT Ok CM about 'You CM about r- CM θ ' 3 CM about un CM θ ' 00 CM © Ok CM ABOUT © CM about 0.26 £ 3 0 | © CM ABOUT © CM © © CM © © CM © ίο CM θ ' Expression B about HIM CM o. o cm 3 CM en about en & cm about © y CM en en Ok 3 ^ CC % CC o. o © CM r » Ok CM ABOUT en en Ok CM en Ok c-t © en 3 ABOUT, CC Expression AND 0.0054 © v *> about about θ ' and 0.0040 about 3 § θ ' © S. about cT 0.0040 $ about about θ ’ about 3 θ ’ CM © ABOUT What about «Ck £ ABOUT © 0.0046: ABOUT 0.0040 0.0040 © § © © 3 s © " § δ θ ' Ξ ffl Ok m en t— © c- en about 3 c- en 3 ITl en Ok en Ok CM 3 © OO vk f3 <3 about m r * 3 cc ec © - © © ca that trt < 01 α Q ω here ABOUT X * ł -4 s 2 ABOUT and- α

PL 208 233 B1

Comments 8 at 3 Cd C. S. £ o <X5 sts cont N AT £ rt e S. - -ABOUT S £ (Z) Q steel comparative steel comparative steel comparative and at and 1 2 for CZ5 G steel comparative steel comparative —I-1 the smaller is The shape of the breakthrough and welds point in the testicle £ cd u o -5 in the testicle pose iadrem pose the nucleus in the testicle pose iadrem pose iadrem X _ □ r ~ 4-1 en ΰ Rating welding- properties: K <1.47 X ABOUT X ABOUT ABOUT X ABOUT ABOUT with-\ '57 * £ .about c_> Hardness K ratio of the weld-output steel (K = max hardness, welds / max hardness of the output steel,) Oh • τ 1.20 izh Ό CM 1.14 1.48 CM * 1) The numbers in the shaded boxes are outside the ranges provided for by the invention. * 2) Assessment of local formability: the hole expansion ratio λ> 60% is expressed by the sign O (good). * 3) Assessment of weldability: the case when the ratio K of the hardness of the weld-base steel (K = max. Hardness of the weld / max. Hardness of the output steel <expressed by the sign O (good). 'c £ ji c | ΰ ° *> £ 498 »N o. o r * ") Oh CM ’Τ 305 376 478 407 380 a Λ - g £ 'J? θ ' «SJ £> 3 »<2 .2 Η <λ o O- v> m en 320 278 242 Γ * Ί f * h 5f CM m 356 314 Rating place- in deform- wholeness λ> 60% X X X ABOUT X X X X Ϊ □ u 2? μ e o About S2 cd £ C? t / 5 O. 2 O £ o. o CM r- Ch r ~ i Tl o. o o. o about 'T 24 rf • 4- 03 1 . c ά 2 'g .2 t- -2 u cu 1026 666 964 694 V) CM © 1109 1011 997 Expression C. about Tl 741 756 428 757 r- r- 5T 1915 1429 Expression D CM CC © 0.24 What θ ' 0.21 0.25 m r * h about' j 0.26. 0.26 Expression B τΤ, cn 3.10 2.84 2.77 2.71 3.62 • n r * T Axis o. o CM 5? from and £ * - Expression AND 0.0040 0.0074 0.0040 0.0040 0.0040 0.0040 6900 '0 CM ABOUT ABOUT about © ' c about 3 about\ m > - 76 m 40 Γ- m TT ΖΊ m about Η 1 Code steel S3 X about -and at C + -. OO X

expressed by the sign of O (good).

Claims (7)

A cold-rolled steel sheet, in particular surface treated, having a tensile strength of at least 780MPa, characterized in that it comprises (% by weight); C: 0.05 to 0.09%, Si: 0.4 to 1.3%, Mn: 2.5 to 3.2%, P: 0.001 to 0.05%, N: 0.0005 to 0.006%, Al: 0.005 to 0.1%, Ti: 0.001 to 0.045% i S within the range defined by the following expression (A), S <0.08 x (Ti (%) - 3.43 x N (%)) + 0.004 (A), in which, when the value of the term Ti (%) - 3.43 x N (%) of the expression (A) is negative , the value is considered to be zero, the remainder being Fe and the inevitable impurities, whereby the indicated components satisfy the following expressions (C) and (D): 950 <(Mneq./(C (%) - (Si (%) / 75))) x percentage of bainite area (%) (C), C (%) + (Si (%) / 20) + (Mn (%) / 18) <0.30 (D), where Mneq. is determined by the following expression (B): Mneq. = Mn (%) - 0.29 x Si (%) + 6.24 x C (%) (B), and the microstructure of the steel sheet has a bainite area percentage of at least 7%, the rest being at least one from ferrite, martensite, tempered martensite and residual austenite. 2. Steel sheet, according to claim A composition according to claim 1, characterized in that it contains as additional components at least one of (in% by weight): Nb: 0.001 to 0.04% B: 0.0002 to 0.0015%, i Mo: 0.05 to 0.50%. 3. Steel sheet, according to claim A composition according to claim 1 or 2, characterized in that it comprises 0.0003 to 0.01% by weight of Ca as an additional component. 4. Steel sheet, according to claim from 1 to 3, characterized in that it contains as an additional component from 0.0002 to 0.01% by weight of Mg. 5. Sheet steel, according to claim The composition of 1 to 4, characterized in that it comprises 0.0002 to 0.01 wt.% of rare earth metals as additional components. Steel sheet according to claim characterized in that it comprises, as additional components, from 0.2 to 2.0% by weight of Cu and from 0.05 to 2.0% by weight of Ni. 7. Steel sheet, according to claim characterized in that it is coated with zinc or its alloy in surface treatment.